CN109477167B - Copper-nickel-tin alloy, method for the production thereof and use thereof - Google Patents

Copper-nickel-tin alloy, method for the production thereof and use thereof Download PDF

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CN109477167B
CN109477167B CN201780044578.9A CN201780044578A CN109477167B CN 109477167 B CN109477167 B CN 109477167B CN 201780044578 A CN201780044578 A CN 201780044578A CN 109477167 B CN109477167 B CN 109477167B
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copper
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nickel
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CN109477167A (en
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凯·韦伯
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Wieland Walker Open Co ltd
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C9/00—Alloys based on copper
    • C22C9/02—Alloys based on copper with tin as the next major constituent
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/001—Continuous casting of metals, i.e. casting in indefinite lengths of specific alloys
    • B22D11/004—Copper alloys
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/10—Alloys containing non-metals
    • C22C1/1036—Alloys containing non-metals starting from a melt
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/10—Alloys containing non-metals
    • C22C1/1036—Alloys containing non-metals starting from a melt
    • C22C1/1073—Infiltration or casting under mechanical pressure, e.g. squeeze casting
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C32/00—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ
    • C22C32/0047—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ with carbides, nitrides, borides or silicides as the main non-metallic constituents
    • C22C32/0052—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ with carbides, nitrides, borides or silicides as the main non-metallic constituents only carbides
    • C22C32/0057—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ with carbides, nitrides, borides or silicides as the main non-metallic constituents only carbides based on B4C
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C32/00—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ
    • C22C32/0047—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ with carbides, nitrides, borides or silicides as the main non-metallic constituents
    • C22C32/0073—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ with carbides, nitrides, borides or silicides as the main non-metallic constituents only borides
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C9/00—Alloys based on copper
    • C22C9/06—Alloys based on copper with nickel or cobalt as the next major constituent
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/08—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of copper or alloys based thereon

Abstract

The invention relates to a high-strength copper-nickel-tin alloy having excellent castability, hot workability and cold workability, high abrasion resistance, high galling resistance and high fretting resistance and improved corrosion resistance and stress relaxation resistance, consisting of, in% by weight: 2.0-10.0% Ni, 2.0-10.0% Sn, 0.01-1.5% Si, 0.002-0.45% B, 0.001-0.09% P, optionally up to a maximum of 2.0% Co, further optionally up to a maximum of 2.0% Zn, optionally up to a maximum of 0.25% Pb, the balance being copper and unavoidable impurities, characterized in that the Si/B ratio of the element content in% by weight of elemental silicon and boron is 0.4 at a minimum and 8 at a maximum; so that the copper-nickel-tin alloy has phases containing Si phase and B phase and Ni-Si-B, Ni-B, Ni-P and Ni-Si system, which obviously improves the processing performance and the service performance of the alloy. The invention also relates to casting variants and further processing variants of high-strength copper-nickel-tin alloys, to a method for producing the same, and to the use of the same.

Description

Copper-nickel-tin alloy, method for the production thereof and use thereof
The invention relates to a copper-nickel-tin alloy having excellent castability, hot formability and cold formability, high resistance to abrasive wear, high resistance to adhesive wear and fretting, and improved corrosion resistance and stress relaxation resistance, according to the preamble of one of claims 1 to 3, a method for the production thereof according to the preamble of claims 9 to 10, and the use thereof according to the preamble of claims 16 to 18.
Due to their good strength properties and their good corrosion resistance and conductivity to heat and current, binary copper/tin alloys are of great importance in both mechanical engineering and motor vehicle construction, as well as in most electronic and electrical engineering.
This group of materials has a high resistance to abrasive wear. Moreover, copper/tin alloys ensure good sliding properties and high fatigue resistance, which makes them extremely suitable for sliding elements in engine and motor vehicle construction and in general mechanical engineering.
Copper-nickel-tin alloys have improved mechanical properties, such as hardness, tensile strength, and yield point, compared to binary copper/tin materials. In this context, the increase in mechanical index is achieved via the hardenability of the Cu-Ni-Sn alloy.
The precipitation process is essential to establish the properties of this group of materials, except that the ratio of the elements nickel and tin is important for the temperature at which spontaneous metastable line (spinodal) segregation occurs in Cu-Ni-Sn alloys.
In the literature, the presence of discontinuous precipitates, in particular at the grain boundaries of the microstructure of Cu-Ni-Sn alloys, is associated with a deterioration of the toughness properties under dynamic stress.
For example, publication DE0833954T1 proposes the production of a metastable-wire Cu-Ni-Sn continuous casting alloy with 8 to 16% by weight Ni, 5 to 8% by weight Sn, and optionally with up to 0.3% by weight Mn, up to 0.3% by weight B, up to 0.3% by weight Zr, up to 0.3% by weight Fe, up to 0.3% by weight Nb, and up to 0.3% by weight Mg, without any processing by kneading. After the solution annealing treatment in the as-cast state and after spinodal aging, the alloy must be rapidly cooled by water quenching in each case in order to obtain a spinodal-segregated microstructure free of discontinuous precipitates.
In contrast, publication DE2350389C states that, with respect to Cu-Ni-Sn alloys having 2 to 98% by weight Ni and 2 to 20% by weight Sn, cold forming with a formation of 75% of at least one degree is necessary in order to be able to prevent the occurrence of embrittlement discontinuity precipitates during the aging annealing.
Document DE69105805T2 mentions the difficulties arising in the mass production of semi-finished products and components from the copper-nickel-tin alloy industry. For example, the occurrence of Sn-rich segregation, particularly at the grain boundaries of the cast microstructure, greatly limits the opportunity for economical further processing. The rich Sn segregation, which cannot be easily eliminated even by thermo-mechanical working operations on Cu-Ni-Sn alloys in the as-cast state, prevents homogeneous distribution of the alloying elements in the matrix. However, this is a fundamental prerequisite for the hardenability of this group of materials. Therefore, it is proposed to finely atomize a melt of a copper alloy with 4 to 18% by weight of Ni and 3 to 13% by weight of Sn, and to collect the spray particles on a collecting surface. The subsequent rapid cooling is intended to counteract the formation of Sn-rich grain boundary segregation.
Document DE4126079C2 discloses that many copper alloys can be produced by conventional block casting methods, followed by hot forming and cold forming, with intermediate annealing operations being only economically less feasible, if at all, because hot forming is difficult due to grain boundary precipitates, segregation or other inhomogeneities.
These copper alloys also include copper-nickel-tin materials. Therefore, in order to ensure cold forming of the alloy in such a cast state, it is recommended to use a thin strip casting method in which the solidification rate of the melt is precisely controlled.
As operating temperatures and pressures rise in modern engines, machines, equipment and units, various different damage mechanisms to individual system components arise. Therefore, from a material and structural point of view, it is necessary to consider not only the type of sliding wear but also the mechanism of damage by vibration friction wear, particularly in the case of designing sliding elements and plug connectors.
Vibratory frictional wear, also known in jargon as fretting, is a type of frictional wear that occurs between vibrating contact surfaces. In addition to the geometric or volumetric wear of the components, the reaction with the surrounding medium leads to fretting corrosion. Damage to the material can significantly reduce the local strength, especially the fatigue strength, of the wear zone. Fatigue cracks may travel from the damaged component surface, and these lead to fatigue fracture/fatigue failure. Under fretting corrosion, the fatigue strength of the component can be reduced to just below the fatigue index of the material.
In one sense, the mechanism of the vibration friction wear is very different from the type of sliding wear in terms of motion. More specifically, in the case of vibration friction wear, the influence of corrosion is particularly significant.
Document DE102012105089a1 describes the consequences of damage caused by the oscillating frictional wear of the plain bearing. To ensure stable positioning of the sliding bearings, they are retracted into the bearing seats. The retracting operation creates significant stresses on the sliding bearing, which stresses are even further increased by increased stresses, thermal expansion and dynamic shaft loads in modern engines. Due to excessive stress, the geometry of the sliding bearing may change, which reduces the original bearing overlap. This enables the sliding bearing to be jogged relative to the bearing seat. These cyclic relative movements with low vibration width at the contact surface between the bearing and the bearing seat lead to vibrational frictional wear/fretting wear of the backing of the sliding bearing. The result is the initiation of cracks and ultimately frictional fatigue failure of the sliding bearing. The results of fretting tests with various sliding bearing materials show that, in particular, Cu-Ni-Sn alloys with Ni contents above 2% by weight, such as the case of metastable wire-hardened copper-nickel-tin alloys, have an insufficient resistance to fretting.
In engines and machines, electrical plug connectors are typically provided in environments where they are subjected to mechanical vibrations. If the elements of the connecting structure are present in different assemblies which, as a result of mechanical stress, are subjected to relative movement with respect to one another, the result may be a corresponding relative movement of the connecting elements. These relative movements lead to vibrational frictional wear and fretting corrosion of the contact areas of the plug connector. Microcracks form in this contact area, which greatly reduces the fatigue resistance of the plug connector material. Failure of the plug connector due to fatigue failure can be a consequence. Furthermore, there is an increase in contact resistance due to fretting corrosion.
Thus, a key factor sufficient to resist vibrational frictional wear/fretting is the combination of material properties of wear resistance, ductility and corrosion resistance.
In order to improve the wear resistance of copper-nickel-tin alloys, a suitable wear matrix must be added to these materials. These wear matrices in the form of hard particles are intended to take on the function of preventing the consequences of abrasive wear and adhesive wear. Hard particles useful in Cu-Ni-Sn alloys include various forms of precipitation.
Document US6379478B1 discloses a teaching of a copper alloy for a plug connector having 0.4 to 3.0% by weight of Ni, 1 to 11% by weight of Sn, 0.1 to 1% by weight of Si and 0.01 to 0.06% by weight of P. It is said that fine precipitates of nickel silicide and nickel phosphide ensure high strength and good stress relaxation resistance of the alloy.
For producing sliding layers on steel-based substrates, document US2129197A names a copper alloy which is applied to a base substrate by applying welding and contains 77 to 92% by weight of Cu, 8 to 18% by weight of Sn, 1 to 5% by weight of Ni, 0.5 to 3% by weight of Si and 0.25 to 1% by weight of Fe. The wear substrate used herein is said to be a silicide and phosphide of the alloying elements nickel and iron.
Document US3392017A discloses a low melting point copper alloy having up to 0.4% by weight of Si, 1 to 10% by weight of Ni, 0.02 to 0.5% by weight of B, 0.1 to 1% by weight of P and 4 to 25% by weight of Sn. The alloy may be applied to a suitable metal substrate surface as a filler material in the form of a cast rod. The alloy has improved ductility compared to the prior art and is machine processable. In addition to being used for deposition soldering, the Cu-Sn-Ni-Si-P-B alloy can be used for deposition by a spray coating method. It is said herein that the addition of phosphorus, silicon and boron is for improving the spontaneous fluidity of the molten alloy and the wettability of the substrate surface, and it is not necessary to use any additional flux.
The teaching disclosed in this document provides for a mandatory Si content of 0.05 to 0.15% by weight and a particularly high P content of 0.2 to 0.6% by weight in the alloy. This emphasizes the main requirement for the spontaneous flow properties of the material. With such a high P content, the hot formability of the alloy will be poor and the metastable line segregability of the microstructure will be insufficient.
According to document US4818307A, the size of the hard particles precipitated in the copper-based alloy has a great influence on its wear resistance. For example, complex silicide/boride formation of elemental nickel and iron up to 5 to 100 μm in size significantly improves the wear resistance of copper alloys having 5 to 30 wt% Ni, 1 to 5 wt% Si, 0.5 to 3 wt% B, and 4 to 30 wt% Fe. Elemental tin is not present in the material. The material is applied as a wear resistant layer to a suitable substrate by deposition welding.
Document US5004581A describes a copper alloy with an additional content of tin in the range of 5 to 15% by weight and/or an additional content of zinc in the range of 3 to 30% by weight, as in the above-mentioned US 4818307A. The addition of Sn and/or zinc improves the resistance of the material to adhesive wear in particular. The material is likewise applied as a wear-resistant layer to a suitable substrate by deposition welding.
However, the copper alloys according to documents US4818307A and US5004581A will have only very limited cold formability due to the required dimensions of 5 μm to 100 μm for the silicide formation/boride formation of the nickel and iron elements required.
The disclosure of precipitation hardenable copper-nickel-tin alloys comes from document US 5041176A. The copper base alloy contains 0.1 to 10% by weight of Ni, 0.1 to 10% by weight of Sn, 0.05 to 5% by weight of Si, 0.01 to 5% by weight of Fe, and 0.0001 to 1% by weight of boron. The material has a content of dispersed intermetallic phases of the Ni-Si system. The properties of the alloy are also illustrated by working examples without any Fe content.
The document KR1020020008710A (abstract) states that metastable wire Cu-Ni-Sn alloys with Sn contents of more than 6% by weight are not thermoformable. The reason is given by Sn-rich segregation at the grain boundaries of the cast microstructure of the Cu-Ni-Sn alloy. Thus, for the disclosed Cu-Ni-Sn multi-species alloys, for high strength wire and sheet, 1 to 8 wt% Ni, 2 to 6 wt% Sn, and 0.1 to 5 wt% of a composition of two or more elements from the Al, Si, Sr, Ti, and B groups are specified.
Document US5028282A discloses a copper alloy with 6 to 25 wt% Ni, 4 to 9 wt% Sn and a further addition in an amount of 0.04 to 5 wt% (alone or together). These further% additions by weight are:
0.03 to 4 percent of Zn, 0.01 to 0.2 percent of Zr,
0.03 to 1.5% of Mn, 0.03 to 0.7% of Fe,
0.03 to 0.5 percent of Mg, 0.01 to 0.5 percent of P,
0.03 to 0.7 percent of Ti, 0.001 to 0.1 percent of B,
0.03 to 0.7% of Cr and 0.01 to 0.5% of Co.
It is stated that the alloying elements Zn, Mn, Mg, P and B are added for deoxidation of the alloy melt. The elements Ti, Cr, Zr, Fe and Co have the functions of grain refinement and strength enhancement.
By alloying with metalloids (e.g. boron, silicon and phosphorus) a relatively high reduction of the base melting temperature can be achieved, which is important for processing purposes. These alloy additions are therefore particularly useful in the field of wear-resistant coating materials and high-temperature materials, including, for example, alloys of the Ni-Si-B system and of the Ni-Cr-Si-B system. Among these materials, in particular, the alloying elements boron and silicon are believed to be responsible for the significant reduction in melting temperature of nickel-based cemented carbides, which makes them useful as self-flowing nickel-based cemented carbides.
The published specification DE2033744B contains important statements about the further function of the alloying element boron in Si-containing metal melts. Hereby, the addition of boron results in a digestion of the oxides formed in the melt and the formation of borosilicate, which rises to the surface of the coating and thus prevents further entry of oxygen. In this way, a smooth surface of the coating can be achieved.
Document DE10208635B4 describes a process in which diffused solder sites of intermetallic phases are present. By diffusion welding, parts having different coefficients of thermal expansion will be joined to each other. In the presence of thermomechanical stresses at the welding site or in the welding operation itself, large stresses are generated at the interface, which can lead to cracks, in particular in the intermetallic phase environment. A remedy proposed is to mix the solder component with particles which achieve a balance of different coefficients of expansion of the connection partners. For example, borosilicate or phosphosilicate particles may minimize thermomechanical stresses in solder bonding due to their favorable coefficient of thermal expansion. Furthermore, the propagation of already induced cracks is hindered by these particles.
The published specification DE2440010B highlights, inter alia, the effect of elemental boron on the electrical conductivity of cast silicon alloys having 0.1 to 2.0% by weight boron and 4 to 14% by weight iron. In this Si-based alloy, a high melting point Si — B phase precipitates, which is called silicon boride.
SiB, generally determined by boron content3、SiB4、SiB6And/or the silicon borides present in SiBn polymorphs differ significantly from silicon in their properties. These silicon borides have metallic properties and are therefore electrically conductive. They have extremely high thermal and oxidative stability. SiB for sintered products is preferred due to its very high hardness and high resistance to abrasive wear6Polymorphs are useful, for example, in ceramic manufacture and ceramic processing.
Conventional Wear-Resistant cemented carbides for surface coatings consist of a relatively ductile matrix of metallic iron, cobalt and Nickel with intergrown silicides and borides as Hard particles (Knotek, O.; Lugchoider, E.; Reimann, H.: Ein Beitrag zur Berteiilung verse Nickel-Boron-Silicon-Hartlegie.A restriction to the Assessment of Wear-Resistant carbide-Boron-Silicon Alloys, Zeitschrift ü r Werkstofftechnik 8(1977)10, p. 331-335). The widespread use of cemented carbides of the Ni-Cr-Si system, Ni-Cr-B system, Ni-B-Si system and Ni-Cr-B-Si system is based on an increase in the wear resistance of these hard particles. Except silicide Ni3Si and Ni5Si2In addition, the Ni-B-Si alloy further contains a boride Ni3B and Ni-Si boride/Ni silicon boride Ni6Si2B. Some slowness of silicide formation in the presence of elemental boron has also been reported. Further investigation of the Ni-B-Si alloy system led to the production of high melting Ni-Si boride Ni6Si2B and Ni4.29Si2B1.43(Lugscheider, E.; Reimann, H.; Knotek, O.: Das Dreisstoffsystem Nickel-Bor-Silicon System, The triphase Nickel-Boron-Silicon System, Monatshefter fur Chemie 106(1975)5, p.1155-1165). These high melting point Ni-Si borides have a relatively broad homogeneity in the direction of boron and silicon.
In many applications, elemental zinc is added to copper-nickel-tin alloys to reduce metal costs. Functionally, the effect of the alloying element zinc is to more significantly form Sn-rich or Ni-Sn-rich phases from the melt. In addition, zinc enhances the formation of precipitates in spinodal Cu-Ni-Sn alloys.
In addition, in many applications, a certain Pb content is also added to the copper-nickel-tin alloy to improve the dry running performance and better material removal workability.
The object of the invention is to provide a high-strength copper-nickel-tin alloy which has excellent hot formability both in the entire nickel content and in the range of tin contents of 2 to 10% by weight in each case. Precursor materials produced by conventional casting methods without the need for spray compaction or strip casting should be useful for thermoforming.
After casting, the copper-nickel-tin alloy should be free of porosity, shrinkage porosity and stress cracks and characterized by a microstructure having a homogeneous distribution of tin-rich phase constituents. Furthermore, intermetallic phases are already present in the microstructure of the copper-nickel-tin alloy after casting. This is important so that the alloy has high strength, high hardness and sufficient wear resistance even in the as-cast condition. Furthermore, it should have high corrosion resistance even in the as-cast state.
First, the as-cast condition of the copper-nickel-tin alloy should not first be homogenized by a suitable annealing treatment in order to be able to establish sufficient hot formability.
With regard to the workability of copper-nickel-tin alloys, the first aim is that their cold formability does not deteriorate significantly, despite the content of intermetallic phases relative to conventional Cu-Ni-Sn alloys. On the other hand, in the case of alloys, the requirement for a minimum degree of forming in the cold forming operation to be carried out should be eliminated. This is considered to be a prerequisite according to the prior art in order to be able to ensure metastable line segregation of the microstructure of the Cu-Ni-Sn material without formation of discontinuous precipitates.
A further requirement for further processing of Cu-Ni-Sn materials corresponding to the prior art is based on the cooling rate after age hardening of the material. Therefore, after spinodal age hardening, it is believed that the material must be rapidly cooled by water quenching to obtain a spinodal line segregated microstructure without discontinuous precipitates. However, as a result of this cooling method dangerous internal stresses may develop after age hardening, it is a further object of the invention to prevent the formation of discontinuous precipitates throughout the manufacturing process, including age hardening, even for alloys.
By further processing operations comprising at least one annealing operation or at least one hot and/or cold forming operation and at least one annealing operation, a microstructure containing fine-grained hard particles can be established which has high strength, high thermal resistance, high hardness, high resistance to stress relaxation and corrosion, sufficient electrical conductivity and high resistance to frictional and vibratory frictional wear mechanisms.
The invention is reflected in respect of copper-nickel-tin alloys by the features of any one of claims 1 to 3, in respect of production processes carried out by the features of claims 9 to 10 and in respect of uses carried out by the features of claims 16 to 18. The other dependent claims relate to advantageous forms and developments of the invention.
The present invention includes a high strength copper-nickel-tin alloy having excellent castability, hot formability and cold formability, high abrasive wear resistance, high adhesive wear resistance and high fretting wear resistance, and improved corrosion resistance and stress relaxation resistance, consisting of, in weight%:
2.0% to 10.0% of Ni,
2.0 to 10.0% of Sn,
0.01 to 1.5% of Si,
0.002% to 0.45% of B,
0.001% to 0.09% of P,
alternatively, up to a maximum of 2.0% Co,
alternatively, up to a maximum of 2.0% Zn,
alternatively, up to a maximum of 0.25% Pb,
and the balance: copper and unavoidable impurities in the form of copper,
it is characterized in that the preparation method is characterized in that,
-the Si/B ratio of the element content in% by weight of the elements silicon and boron has a minimum value of 0.4 and a maximum value of 8;
the copper-nickel-tin alloy comprises phases containing Si, B and the system Ni-Si-B, Ni-B, Ni-P and Ni-Si, which significantly improve the workability and usability of the alloy.
The present invention also includes a high strength copper-nickel-tin alloy having excellent castability, hot formability and cold formability, high abrasive wear resistance, high adhesive wear resistance and high fretting wear resistance, and improved corrosion resistance and stress relaxation resistance, consisting of, in weight percent:
2.0% to 10.0% of Ni,
2.0 to 10.0% of Sn,
0.01 to 1.5% of Si,
0.002% to 0.45% of B,
0.001% to 0.09% of P,
alternatively, up to a maximum of 2.0% Co,
alternatively, up to a maximum of 2.0% Zn,
alternatively, up to a maximum of 0.25% Pb,
and the balance: copper and unavoidable impurities in the form of copper,
it is characterized in that the preparation method is characterized in that,
-the Si/B ratio of the element content in% by weight of the elements silicon and boron has a minimum value of 0.4 and a maximum value of 8;
the following microstructural elements are present in the alloy after casting:
a) a Si-containing and P-containing metal-based component having, based on the overall microstructure:
a1) can be calculated by empirical formula CuhNikSnmReported, and having a first phase composition of up to 35% by volume of the (h + k)/m ratio of the element content in 2 to 6 atomic%,
a2) can be calculated by empirical formula CupNirSnsReported, and having a second phase composition of up to 15% by volume of the (p + r)/s ratio of the element content in 10 to 15 atomic%,
a3) the balance of the solid copper solution;
b) based on the phases present in the overall microstructure,
b1) as a Si-containing phase and a B-containing phase in an amount of 0.01 to 10% by volume,
b2) based on 1% by volume to 15% by volume, as Ni having the empirical formulaxSi2B (x is 4 to 6) Ni-Si boride,
b3) based on 1 to 15% by volume, as Ni boride,
b4) as Ni phosphide in an amount of 1 to 5% by volume,
b5) based on 1 to 5 volume percent as Ni silicide
In the microstructure, they are present individually and/or as addition compounds and/or mixture compounds and are encapsulated by tin and/or the first phase component and/or the second phase component;
during casting, the Si-containing and B-containing phases, Ni-Si boride and Ni boride, Ni phosphide and Ni silicide, present alone and/or as silicon boride in addition compounds and/or mixed compounds, constitute seeds which crystallize homogeneously during solidification/cooling of the melt, so that the first phase constituents and/or the second phase constituents are distributed homogeneously in the microstructure, for example islands-like and/or network-like;
si-containing and B-containing phases in the form of borosilicates and/or borophosphosilicates act together with phosphosilicates as wear-and corrosion-protective coatings on alloyed semi-finished materials and components.
Advantageously, the first phase component and/or the second phase component is present in the cast microstructure of the alloy in at least 1% by volume.
The homogeneous distribution of the first phase component and/or the second phase component in the form of islands and/or networks means that the microstructure is not segregated. This segregation is understood to mean the accumulation of the first phase component and/or the second phase component in the cast microstructure, which takes the form of grain boundary segregation, which, under thermal and/or mechanical stress at the time of casting, can lead to damage to the microstructure, which is a form of cracks that can lead to fractures. The as-cast microstructure is still free of pores, shrinkage porosity, stress cracks and discontinuous precipitates of the (Cu, Ni) -Sn system.
In this variant, the alloy is in the as-cast condition.
The present invention also includes a high strength copper-nickel-tin alloy having excellent castability, hot formability and cold formability, high abrasive wear resistance, high adhesive wear resistance and high fretting wear resistance, and improved corrosion resistance and stress relaxation resistance, consisting of, in weight percent:
2.0% to 10.0% of Ni,
2.0 to 10.0% of Sn,
0.01 to 1.5% of Si,
0.002% to 0.45% of B,
0.001% to 0.09% of P,
alternatively, up to a maximum of 2.0% Co,
alternatively, up to a maximum of 2.0% Zn,
alternatively, up to a maximum of 0.25% Pb,
and the balance: copper and unavoidable impurities in the form of copper,
it is characterized in that the preparation method is characterized in that,
-the Si/B ratio of the element content in% by weight of the elements silicon and boron has a minimum value of 0.4 and a maximum value of 8;
-after further treatment of the alloy by at least one annealing operation or by at least one hot and/or cold forming operation and at least one annealing operation, the following microstructure constituents are present:
A) a metal-based component having the following overall microstructure,
A1) can be calculated by empirical formula CuhNikSnmReported, and having a first phase composition of up to 15% by volume of the (h + k)/m ratio of the element content in 2 to 6 atomic%,
A2) can be calculated by empirical formula CupNirSnsReported, and having a second phase composition of up to 5% by volume of the (p + r)/s ratio of the element content in 10 to 15 atomic%, and
A3) the balance of the solid copper solution;
B) based on the overall microstructure, the phases are present in such a way,
B1) in the microstructure as Si-containing phase and B-containing phase in an amount of 2 to 30% by volume, having the empirical formula NixSi2Ni-Si borides of B (x ═ 4 to 6) as Ni borides, Ni phosphides and Ni silicides, present individually and/or as addition compounds and/or mixed compounds, and coated with precipitates of the (Cu, Ni) -Sn system,
B2) as a continuous precipitate of the (Cu, Ni) -Sn system in the microstructure in amounts of up to 80% by volume,
B3) as Ni phosphide and Ni silicide, present alone and/or as addition compounds and/or mixed compounds, encapsulated by precipitates of the (Cu, Ni) -Sn system and having a size of less than 3 μm, in a microstructure, in a volume of 2 to 30%;
the Si-containing phase and the B-containing phase, Ni-Si boride and Ni boride, Ni phosphide and Ni silicide in the form of silicon boride, present alone and/or as addition compounds and/or mixed compounds, constitute seeds for static and dynamic recrystallization of the microstructure during further processing of the alloy, which enables the establishment of a homogeneous and fine-grained microstructure;
the Si-containing phase and the B-containing phase in the form of borosilicate and/or the borophosphosilicate together with the phosphosilicate act as a wear-and corrosion-protective coating on the semifinished materials and components of the alloy.
Advantageously, the continuous precipitates of the (Cu, Ni) -Sn system are present in the microstructure of the alloy in the further processed state in at least 0.1% by volume.
The microstructure is not segregated even after further processing of the alloy. This segregation is understood to mean an accumulation of the first phase component and/or the second phase component in the microstructure, which takes the form of grain boundary segregation, which, particularly under dynamic stresses on the component, can lead to damage to the microstructure, which is in the form of cracks that can lead to fractures.
The microstructure of the alloy after further processing is free of air holes, shrinkage cavities and stress cracks. As an essential feature of the invention, it should be emphasized that the microstructure in the further processing state is free of discontinuous precipitates of the (Cu, Ni) -Sn system.
In this second variant, the alloy is in a further processed state.
The present invention has been made in view of providing a copper-nickel-tin alloy having a Si-containing phase and a B-containing phase and a phase having a Ni-Si-B system, a Ni-P system and a Ni-Si system. These phases significantly improve the processability in castability, hot formability and cold formability. In addition, these phases improve the performance properties of the alloy by increasing strength and resistance to abrasive wear, adhesive wear and fretting. These phases additionally improve corrosion resistance and stress relaxation resistance as further performance properties of the invention.
The copper-nickel-tin alloy of the invention can be produced by a sand casting process, a shell mould casting process, a precision casting process, a full mould casting process, a pressure die casting process, a lost foam process and a permanent mould casting process or with the help of a continuous or semi-continuous casting process.
It is possible to use a primary forming technique which is complicated in terms of processing technique and cost, but is not absolutely necessary for the production of the copper-nickel-tin alloy of the invention. For example, the use of spray compaction or thin strip casting may be omitted. The cast form of the copper-nickel-tin alloy according to the invention can be hot-formed directly, for example by hot rolling, continuous casting or forging, in particular over the entire range of Sn and Ni contents, without it being absolutely necessary to carry out a homogenizing anneal. It is also worth noting that after the shell mold or continuous casting form made from the alloy of the present invention, there is no need to perform any complicated forging or compression process at elevated temperatures to weld (i.e., close) pores and cracks in the material. Thus, the processing-related limitations that have hitherto been present in the production of semi-finished products and components from copper-nickel-tin alloys are further eliminated.
Depending on the casting process, the microstructure of the metal-based material of the copper-nickel-tin alloy of the invention in the as-cast state consists of an increased proportion of tin-rich phases (alpha phases) uniformly distributed in the solid copper solution, with an increase in the Sn content of the alloy.
The tin-rich phase of these metal-based materials can be divided into a first phase component and a second phase component. The first phase composition may be determined by empirical formula CuhNikSnmReported, and has a (h + k)/m ratio of the element content in 2 to 6 atomic%. The second phase composition may be determined by empirical formula CupNirSnsReported, and has a (p + r)/s ratio of the element content in 10 to 15 atomic%.
The alloy of the present invention is characterized by a Si-containing phase and a B-containing phase, which can be divided into two groups.
The first group relates to the silicon-containing phase and the B-containing phase, which are in the form of silicon borides, and SiB may be present3、SiB4、SiB6And SiBnAmong the polymorphs. "n" in the compound SiBn represents high solubility of elemental boron in the silicon lattice.
The second group of Si-containing and B-containing phases relates to silicate compounds of borosilicate and/or borophosphosilicate.
In the copper-nickel-tin alloy of the present invention, the microstructure components of the Si-containing phase and the B-containing phase in the form of silicon boride and in the form of borosilicate and/or borophosphosilicate are not less than 0.01% by volume and not more than 10% by volume.
The homogeneous arrangement of the constituents of the first phase and/or of the second phase in the microstructure of the alloy according to the invention is brought about in particular by the action of a silicon-containing phase and a boron-containing phase in the form of silicon boride and Ni-Si boride withEmpirical formula NixSi2B (x ═ 4 to 6), predominantly has precipitated out of the melt. Subsequently, during solidification/cooling of the melt, Ni borides are preferably precipitated on the already present silicon boride and Ni-Si borides. The boron compound, which is present alone and/or as an addition compound and/or a mixture compound, in its entirety, serves as primary seed during the first solidification/cooling of the melt.
Later on in the solidification/cooling of the melt, the Ni phosphide and Ni silicide precipitate preferentially as secondary seeds on the primary seeds of silicon boride, Ni-Si boride and Ni boride, which have been present individually and/or as addition compounds and/or mixed compounds.
The Ni-Si boride and the Ni boride are each present in the microstructure at 1% volume to 15% volume, and the Ni phosphide and the Ni silicide are each present at 1% volume to 5% volume.
Thus, in the microstructure, the Si-containing phase and the B-containing phase in the form of silicon boride, have the empirical formula NixSi2The Ni — Si borides and Ni borides, Ni phosphides and Ni silicides of B (x ═ 4 to 6) are present individually and/or as addition compounds and/or mixed compounds.
These phases are hereinafter referred to as crystallization seeds.
Finally, the elemental tin and/or the first phase component and/or the second phase component of the metal-based material preferably crystallize in the region of the crystallization seeds, as a result of which the crystallization seeds of the tin and/or the first phase component and/or the second phase component are enveloped.
These crystalline seeds, which are encapsulated by tin and/or the first phase component and/or the second phase component, are hereinafter referred to as hard particles of the first type.
In the as-cast condition of the alloy of the present invention, the first hard particles are less than 80 μm in size. Advantageously, the first type of hard particles has a size of less than 50 μm.
As the Sn content of the alloy increases, the arrangement of the first phase component and/or the second phase component in the form of islands is converted into a network arrangement in the microstructure.
In the cast microstructure of the copper-nickel-tin alloy of the present invention, the first phase component may account for up to 35% by volume. The second phase component comprises a microstructural fraction of up to 15% by volume. Advantageously, the first phase component and/or the second phase component is present in at least 1% by volume in the as-cast microstructure of the alloy.
Due to the addition of the alloying element boron, there are phosphides and suicides which are suppressed and thus are only incompletely formed during casting of the alloys of the invention. For this reason, the contents of phosphorus and silicon remain dissolved in the metal-based material in the as-cast state.
Conventional copper-nickel-tin alloys have a relatively wide solidification interval. Such wide solidification intervals in casting increase the risk of gas absorption and thus incomplete, coarse, often dendritic crystallization of the melt. The result is usually porosity and coarse tin enrichment segregation, at whose phase boundaries shrinkage porosity and stress cracking often occur. In this group of materials, Sn-rich segregation also occurs preferentially at grain boundaries.
By virtue of the combined contents of boron, silicon and phosphorus, various processes in the alloy melt of the invention are activated, which critically alter its solidification characteristics compared to conventional copper-nickel-tin alloys.
In the melt according to the invention, the elements boron, silicon and phosphorus have a deoxidizing function. The addition of boron and silicon makes it possible to reduce the phosphorus content without reducing the deoxidation strength of the melt. With this measure, the adverse effect of sufficient deoxidation of the melt can be suppressed by the addition of phosphorus. A high P content will therefore additionally extend the solidification interval of the copper-nickel-tin alloy, which is already in any case very large, which will lead to an increased tendency to pores and segregation in this material type. The adverse effects of phosphorus are reduced by limiting the P content in the alloys of the present invention to the range of 0.001 to 0.09% by weight.
In particular, lowering the base melting temperature by elemental boron and crystallization seeds results in a reduction in the solidification interval of the alloys of the present invention. As a result, the as-cast condition of the present invention has a very uniform microstructure with a fine distribution of the respective phase components. Therefore, in the alloy of the present invention, tin-rich segregation does not occur particularly at grain boundaries.
In the melt of the alloy of the invention, the effect of the elements boron, silicon and phosphorus is the reduction of the metal oxide. The elements themselves are simultaneously oxidized and generally rise to the surface of the casting, where they form a protective layer in the form of borosilicate and/or borophosphosilicate and phosphosilicate, which protects the casting from gas absorption. A particularly smooth surface of the alloy castings of the present invention was found, indicating the formation of such a protective layer. The as-cast microstructure of the present invention also has no porosity throughout the cross-section of the casting.
In the context of comments relating to the cited documents, the advantage of introducing borosilicate and phosphosilicate to avoid stress cracks between phases having different coefficients of thermal expansion during diffusion welding is mentioned.
The basic concept of the invention is to apply the effects of borosilicate, borophosphosilicate and phosphosilicate in respect of matching the different coefficients of thermal expansion of the connection partners in diffusion welding with the casting, thermoforming and heat treatment processes of cupronickel-tin materials. Due to the wide solidification interval of these alloys, high mechanical stresses occur between the low Sn and Sn-rich structural regions, which crystallize in an offset manner and can lead to cracks and pores. In addition, these damage characteristics can also occur during the hot forming and high temperature annealing operations of copper-nickel-tin alloys due to different hot forming characteristics and different coefficients of thermal expansion for the low Sn and Sn-rich microstructure constituents.
Effect of the combined addition of boron, silicon and phosphorus on the copper-nickel-tin alloy according to the invention during solidification of the melt, the microstructure of the first phase component and/or the second phase component of the metal-based material, which is distributed homogeneously, is first of all in the form of islands and/or in the form of networks, by the action of crystallization seeds. In addition to the crystallization seeds, the formation of the Si-containing phase and the B-containing phase together with the phosphosilicate in the form of borosilicate and/or borophosphosilicate during the solidification of the melt ensures the necessary matching of the coefficients of thermal expansion of the first phase component and/or the second phase component and the solid copper solution of the metal-based material. In this way, pores and stress cracks are prevented from forming between phases having different Sn contents.
A further effect of the alloy content of the copper-nickel-tin alloy of the invention is a significant change in the as-cast grain structure. Thus, it was found that in the primary cast microstructure, a substructure with a grain size of sub-grains smaller than 30 μm was formed.
Alternatively, the alloy of the present invention may be further processed by annealing or by hot and/or cold forming operations and at least one annealing operation.
One way to further process the copper-nickel-tin alloy of the present invention is to convert the casting into a final form with the desired properties by at least one cold forming operation and at least one annealing operation.
The alloy of the present invention has high strength even in the as-cast condition due to the uniform cast microstructure and the first type of hard particles precipitated therein. As a result, the castings have relatively low cold formability, which makes it difficult to further economically process them. For this reason, it has been found advantageous to perform a homogenizing annealing operation on the casting prior to the cold forming operation.
To ensure the age hardenability of the present invention, accelerated cooling after the homogenizing annealing process has been found to be advantageous. It has been found here that cooling methods with a relatively low cooling rate can be used in addition to water quenching, due to the slowness of the precipitation mechanism and the separation mechanism. For example, it has been found that the use of accelerated air cooling is equally feasible in order to reduce the effects of hardness enhancement and strength increase of precipitation processes and detachment processes in the microstructure to a sufficient extent during the homogenizing annealing operation of the present invention.
The significant effect of the crystallization seeds on the recrystallization of the microstructure according to the invention is manifested in the microstructure which can be established after cold forming by annealing in a temperature range from 170 ℃ to 880 ℃ and an annealing time between 10 minutes and 6 hours. The particularly fine structure of the recrystallized alloy enables a further cold forming step, the degree of forming of which is typically greater than 70%. In this way, an ultra-high strength state of the alloy can be established.
In the further processing of the inventionThe tensile strength R can be established by these highly cold forming processes which have become possiblemYield point Rp0.2And particularly high values of hardness. Especially Rp0.2The level of the parameter is important for the sliding element and the guiding element. Further, Rp0.2The high value of (a) is a prerequisite for the necessary spring properties of plug connectors in electronics and electrical engineering.
In the description of many documents describing the prior art relating to the processing and performance of copper-nickel-tin materials, reference is made to the need to observe a minimum degree of cold forming of 75%, for example, in order to prevent the precipitation of discontinuous precipitates of the (Cu, Ni) -Sn system in the microstructure.
In contrast, regardless of the degree of cold forming, the microstructure of the alloy of the present invention remains free of discontinuous precipitates of the (Cu, Ni) -Sn system. For example, for a particularly advantageous embodiment of the invention, it was found that the microstructure of the invention is free of discontinuous precipitates of the (Cu, Ni) -Sn system even in the case of very small degrees of cold forming of less than 20%.
According to the prior art, conventional, metastable-linearly segregatable Cu-Ni-Sn materials are considered to be thermoformable, but with great difficulty, if any.
The effect of crystallization seeds is also observed during the process of hot forming the copper-nickel-tin alloy of the present invention. Considering the crystallization seeds is mainly related to the fact that the dynamic recrystallization in the hot forming of the alloy of the invention preferentially takes place in the temperature range of 600 ℃ to 880 ℃. This leads to a further increase in the homogeneity of the microstructure and the fine particle size.
Advantageously, the cooling of the semifinished products and components after thermoforming can be effected with calm or accelerated air or with water.
As in the case after casting, it is also possible to establish a particularly smooth component surface after the hot forming of the casting. This observation indicates that the formation of Si-containing phases and B-containing phases and phosphosilicates in the form of borosilicate and/or borophosphosilicate occurs in the material during thermoforming. Even during thermoforming, the silicate together with the crystallization seeds leads to a matching of the different coefficients of thermal expansion of the phases of the metal-based material of the invention. Thus, the surface and microstructure of the hot formed part (as in the case after casting) is also free of cracks and voids after hot forming.
Advantageously, the at least one annealing treatment of the as-cast and/or as-hot-formed condition of the invention may be carried out at a temperature ranging from 170 ℃ to 880 ℃ for a duration of 10 minutes to 6 hours, and alternatively by cooling under calm or accelerated air or with water.
One aspect of the invention relates to an advantageous process for further processing comprising carrying out at least one cold forming operation in a cast state or in a hot formed state or in an annealed cast state or in an annealed hot formed state.
Preferably, the at least one annealing treatment of the cold formed state of the invention can be carried out at a temperature ranging from 170 ℃ to 880 ℃ for a duration of 10 minutes to 6 hours, and alternatively by cooling under calm or accelerated air or with water.
Advantageously, the stress relief annealing/age hardening annealing operation may be performed at a temperature in the range of 170 ℃ to 550 ℃ for a time period in the range of 0.5 hours to 8 hours.
After further treatment of the alloy by at least one annealing operation or by at least one hot and/or cold forming operation and at least one annealing operation, precipitates of the (Cu, Ni) -Sn system are preferably formed in the region of the crystallization seeds, as a result of which the crystallization seeds are enveloped by these precipitates.
These crystalline seeds, which are surrounded by precipitates of the (Cu, Ni) -Sn system, are hereinafter referred to as hard particles of the second type.
As a result of further processing of the alloy of the invention, the size of the second type of hard particles is reduced compared to the size of the first type of hard particles. The size of the second type of hard particles decreases in particular with increasing degree of cold forming, since these are the hardest components in the alloy and cannot contribute to the shape change of the metal-based material surrounding them. Depending on the degree of cold forming, the resulting hard particles of the second type and/or the resulting segments of the hard particles of the second type have a size of less than 40 μm to even less than 5 μm.
The Ni content and Sn content of the present invention each vary within the range of 2.0 wt% to 10.0 wt%. Ni content and/or Sn content below 2.0 wt.% can result in excessively low strength and hardness values. Moreover, the alloy will have insufficient running properties under sliding stress. The resistance of the alloy to abrasive wear and adhesive wear is not satisfactory. When the Ni content and/or the Sn content is more than 10.0% by weight, the toughness of the alloy of the present invention is rapidly deteriorated, with the result that the dynamic durability of a member made of the material is reduced.
With regard to the assurance of optimum dynamic durability of the components made of the alloy according to the invention, it was found to be advantageous for the content of nickel and tin to be in each case in the range from 3.0% by weight to 9.0% by weight. In this respect, for the invention, a range of 4.0% by weight to 8.0% by weight is particularly preferred for the content of the elements nickel and tin in each case.
As regards copper materials containing Ni and Sn, it is known from the prior art that the spinodal segregation degree of the microstructure increases with increasing Ni/Sn element content ratio of the elements nickel and tin in% by weight. This is true for Ni and Sn contents above and above about 2% by weight. With a decreasing Ni/Sn ratio, the precipitate formation mechanism of the (Cu, Ni) -Sn system gains more weight, which leads to a reduction of the metastable linearly segregated microstructure fraction. One particular result is that, as the Ni/Sn ratio decreases, discrete precipitates of the (Cu, Ni) -Sn system are formed to a greater extent.
The basic features of the copper-nickel-tin alloys of the present invention include a severe suppression of the effect of the Ni/Sn ratio on the formation of discontinuous precipitates in the microstructure. Thus, it has been found that in the microstructure of the present invention there is no precipitation of discrete precipitates of the (Cu, Ni) -Sn system, largely independent of the Ni/Sn ratio and independent of the age hardening conditions.
In contrast, during further processing of the alloy according to the invention, continuous precipitates of the (Cu, Ni) -Sn system are formed at up to 80% by volume. Advantageously, the continuous precipitates of the (Cu, Ni) -Sn system are present in the microstructure of the alloy in the further processed state in at least 0.1% by volume.
The effect of the crystallization seeds during solidification/cooling of the melt, the effect of the crystallization seeds as recrystallization seeds and the effect of the silicate groups with respect to wear protection and corrosion protection only achieve a certain degree of technical significance of the alloy according to the invention when the silicon content is at least 0.01% by weight and the boron content is at least 0.002% by weight. In contrast, if the Si content exceeds 1.5% by weight and/or the B content exceeds 0.45% by weight, this results in deterioration of casting characteristics. Too high a crystalline seed content causes the melt to become thicker and thicker. In addition, the result is a reduction in the toughness of the alloys of the invention.
An advantageous range of Si content is considered to be within the limit of 0.05 wt% to 0.9 wt%. A particularly advantageous content of silicon has been found to be from 0.1% by weight to 0.6% by weight.
For elemental boron, a content of 0.01 to 0.4% by weight is considered advantageous. Boron contents of from 0.02 to 0.3% by weight have been found to be particularly advantageous.
In order to ensure a sufficient content of Ni-Si boride and Si-containing and B-containing phases in the form of borosilicate and/or borophosphosilicate, a lower limit of the elemental ratio of silicon and boron elements has been found to be important. Thus, in the alloy of the present invention, the minimum Si/B ratio of the elemental contents in% by weight of the elements silicon and boron is 0.4. For the alloys of the present invention, the element content in% by weight of the elements silicon and boron has an advantageous minimum Si/B ratio of 0.8. Preferably, the minimum Si/B ratio of the elemental contents in% by weight of the elements silicon and boron is 1.
For another important feature of the invention, it is important that the fixed upper limit of the Si/B ratio of the elemental contents in% by weight of elemental silicon and boron is 8. After casting, a portion of the silicon is dissolved in the metal-based material and incorporated in the first type of hard particles.
During the thermal or thermomechanical further processing operation in the as-cast state, the silicide component of the first type of hard particles is at least partially dissolved. This increases the Si content of the metal-based material. If this exceeds the upper limit, an excessive proportion of Ni silicide may precipitate as the size increases. These will critically reduce the cold formability of the present invention.
For this purpose, the maximum Si/B ratio of the elemental contents in% by weight of the elements silicon and boron of the alloy of the invention is 8. By this measure, the size of the Ni silicide formed during the processing operation in the hot or thermo-mechanical further as cast state of the alloy can be reduced to below 3 μm. In addition, this limits the content of Ni silicide. In this respect, it has been found to be particularly advantageous to limit the Si/B ratio of the elemental contents in% by weight of the elements silicon and boron to a maximum value of 6.
Precipitation of crystallization seeds affects the viscosity of the alloy melt of the present invention. This fact underscores why the addition of phosphorus is essential. The effect of the phosphorus is that the melt has sufficient fluidity despite the presence of the crystallization seeds, which is of great significance for the castability according to the invention. The phosphorus content of the alloy of the invention is from 0.001 to 0.09% by weight.
Below 0.001% by weight, the P content no longer contributes to ensuring sufficient castability according to the invention. If the phosphorus content of the alloy assumes values above 0.09% by weight, on the one hand, excessively large Ni components are incorporated in the form of phosphides, which reduce the spinodal separability of the microstructure. On the other hand, in the case where the P content is more than 0.09% by weight, the thermoformability of the present invention may be seriously deteriorated. For this reason, P contents of from 0.01% by weight to 0.09% by weight have been found to be particularly advantageous. The P content is preferably from 0.02 to 0.08% by weight.
For another reason, the alloying element phosphorus is of great importance. Together with the desired maximum Si/B ratio of the element contents in% by weight of the elements silicon and boron of 8, which can be attributed to the phosphorus content of the alloy, Ni phosphides and Ni suicides (size not more than 3 μm and contents of 2% by volume up to 30% by volume) which are present alone and/or as addition compounds and/or mixed compounds and which are encapsulated by precipitates of the (Cu, Ni) -Sn system can be formed in the microstructure after further processing of the invention.
These Ni phosphides and Ni silicides, which are present individually and/or as addition compounds and/or mixed compounds and are encapsulated by precipitates of the (Cu, Ni) -Sn system, and have a size of not more than 3 μm, are hereinafter referred to as hard particles of the third type.
In the microstructure of the invention in the further processed state of particularly preferred configurations, the hard particles of the third type even have a size of less than 1 μm.
These third type of hard particles first complement the second type of hard particles in their function as wear substrates. Thus, they increase the strength and hardness of the metal-based material, thereby increasing the resistance of the alloy to abrasive wear stresses. Second, the third type of hard particles increases the resistance of the alloy to adhesive wear. Finally, the effect of these third type hard particles is a critical increase in the thermal strength and stress relaxation resistance of the alloys of the present invention. This is an important prerequisite for the alloys according to the invention, in particular for sliding elements and components and connecting elements in electrical/electronic engineering.
The alloy of the present invention has the characteristics of a precipitation-hardenable material due to the content of the first type of hard particles in the as-cast microstructure and the second and third types of hard particles in the as-further-processed microstructure. Advantageously, the invention corresponds to a precipitation-hardenable and metastable-line segregatable copper-nickel-tin alloy.
The sum of the elemental contents of the elements silicon, boron and phosphorus is advantageously at least 0.2% by weight.
The casting variants and further processing variants of the alloys of the invention may comprise the following optional elements:
elemental cobalt may be added to the copper-nickel-tin alloy of the present invention at levels up to 2.0% by weight. Due to the similarity relationship between nickel and cobalt elements and due to the similar Si silicon boride formation, silicide formation and phosphide formation properties of cobalt relative to nickel, the alloying element cobalt may be added to participate in the formation of crystalline seeds and the formation of hard particles of the first, second and third types in the alloy. As a result, the content of Ni incorporated in the hard particles can be reduced. This can achieve the effect of: the Ni content for metastable line segregation of the microstructure that is effectively usable in metal-based materials is increased. The strength and hardness of the present invention can be significantly improved by advantageously adding 0.1 to 2.0% by weight of Co.
Elemental zinc may be added to the copper-nickel-tin alloy of the present invention in an amount of 0.1 to 2.0% by weight. It was found that the alloying element zinc increases the proportion of the first phase component and/or the second phase component in the metal-based material according to the invention, depending on the Ni content and the Sn content of the alloy, which leads to an increase in strength and hardness. The interaction between the Ni component and the Zn component is considered to be the cause of this. Due to these interactions between the Ni component and the Zn component, a reduction in the first type of hard particle size and the second type of hard particle size is also found, which consequently results in a finer distribution in the microstructure.
Below 0.1% by weight of Zn, it is not possible to observe these effects on the microstructure and mechanical properties of the present invention. When the Zn content is more than 2.0% by weight, the toughness of the alloy is lowered to a low level. The corrosion resistance of the copper-nickel-tin alloy of the present invention is also deteriorated. Advantageously, a zinc content in the range of 0.1 to 1.5% by weight may be added to the present invention.
Optionally, a small proportion of lead up to a maximum of 0.25% by weight above the contamination limit may be added to the copper-nickel-tin alloy of the invention. In a particularly preferred advantageous embodiment of the invention, the copper-nickel-tin alloy is free of lead apart from any unavoidable contamination, which complies with current environmental standards. In this regard, lead content is expected to be as high as 0.1% by weight of Pb max.
The formation of Si-containing phases and B-containing phases and phosphosilicates in the form of borosilicate and/or borophosphosilicate not only results in a significant reduction of the porosity and crack content in the microstructure of the alloy of the invention. These silicate-based phases also act as wear protection and corrosion protection coatings on the components.
During the adhesive wear stresses on components made of the copper-nickel-tin alloy according to the invention, the alloying element tin plays a particular role in the formation of the so-called friction layer between the friction partners. This mechanism is important, particularly under mixed friction conditions, as the idle performance of the material becomes more and more important. The friction layer results in a reduction in the size of the pure metal contact area between the friction partners, which prevents welding or fretting of the components.
Even higher operating pressures and operating temperatures occur as a result of the increased efficiency of modern engines, machines and units. This is particularly observed in newly developed internal combustion engines where the goal is more complete combustion of the fuel. In addition to the elevated temperature of the space surrounding the internal combustion engine, there is also an evolution of the heat that occurs during the operation of the plain bearing system. Due to the high temperatures in the operation of the bearing, Si-containing phases and B-containing phases and phosphosilicates in the form of borosilicate and/or borophosphosilicate are formed in the component made of the alloy of the invention, similar to during casting and hot forming. These compounds also enhance the friction layer formed mainly by the alloying element tin, which results in an increase in the resistance to adhesive wear of sliding elements made of the alloy of the invention.
Thus, the alloy of the present invention ensures a combination of wear and corrosion resistance. This combination of properties leads to the required high resistance to the tribological wear mechanism and high material resistance to fretting corrosion. In this way, the present invention has excellent applicability as a sliding member and a plug connector because it has high resistance to sliding wear and resistance to vibration friction wear called fretting.
In addition to the important contribution of the hard particles of the third type to the mechanism of increasing the wear resistance and adhesion resistance of the frictional wear of the present invention, the hard particles of the third type play a crucial role in increasing the vibration resistance. Together with the hard particles of the second type, the hard particles of the third type constitute a barrier to the propagation of fatigue cracks which can be introduced into the stress member, in particular under vibrational frictional wear (known as fretting). Thus, with respect to the increase in the resistance to fretting wear (known as fretting) of the alloys of the invention, the hard particles of the second and third type are particularly supplemented with the wear and corrosion protection of the Si-containing phase and the B-containing phase and phosphosilicate in the form of borosilicate and/or borophosphosilicate.
Heat resistance and stress relaxation resistance are further fundamental properties of alloys for end uses where higher temperatures occur. In order to ensure sufficiently high heat resistance and stress relaxation resistance, a high density of fine precipitates is considered to be advantageous. This precipitate is a continuous precipitate of hard particles of the third type and the (Cu, Ni) -Sn system in the alloy of the present invention.
The alloy of the present invention, even in the as-cast condition, has a high degree of strength, hardness, ductility, composite wear and corrosion resistance due to the uniform and fine-grained microstructure and the content of hard particles of the first type that are substantially free of porosity, cracks and segregation. This combination of properties means that the sliding element and the guide element can even be produced from cast form. The cast state of the invention can additionally be used for producing housings for fittings and for water, oil and fuel pumps. The alloy of the invention can also be used for propellers, wings, screws and hubs for shipbuilding.
Further processed variants of the invention may find use in fields of application with particularly high complex and/or dynamic component stresses.
The excellent strength properties and wear resistance as well as corrosion resistance of the copper-nickel-tin alloy of the invention mean possible further uses. Thus, the present invention is applicable to metal products in structures for the cultivation of marine habitats (aquaculture). Furthermore, the invention can be used for the production of pipes, seals and connecting bolts required in the marine and chemical industries.
This material is of great significance for the use of the alloy of the present invention to produce percussion instruments. High quality cymbals in particular have so far been produced from copper alloys, usually containing tin, usually by thermoforming and at least one annealing operation, before they are usually transformed into the final shape by a bell or a shell. Subsequently, the cymbal is annealed again before the material removal finishing process. Therefore, the production of various variants of cymbals (e.g. rhythm, hi-hat, hanging cymbals, chinese cymbals, cymbals and effect cymbals) requires particularly advantageous thermo-formability of the material, which is ensured by the alloy of the present invention. Within the limits of the range of chemical compositions according to the invention, the different microstructural compositions of the phases of the metal-based material and of the different hard particles can be set within wide limits. In this way, the acoustic properties of the cymbal can be affected even from the perspective of the alloy.
In particular for the production of composite plain bearings, the invention can be used for application to composite counterparts by joining methods. Thus, the composite production between the sheet, plate or strip of the invention and the steel cylinder or strip, preferably made of quenched and tempered steel, can be achieved by forging, brazing or welding, optionally with at least one annealing operation in the temperature range 170 ℃ to 880 ℃. For example, the composite bearing cup or the composite bearing bushing can likewise be produced by roll cladding, induction or conduction roll cladding or by laser roll cladding, optionally also with at least one annealing operation in the temperature range from 170 to 880 ℃.
The formation of microstructures in the alloys of the present invention creates a further option for the production of composite sliding elements, such as composite bearing cups or composite bearing bushes. For example, a coating of tin or a Sn-rich material for use as a running layer in bearing operation can be applied on the substrate of the invention by hot dip or electrolytic tin plating, sputtering or by a PVD method or a CVD method.
In this way, high performance composite sliding elements such as composite bearing cups or composite bearing bushes can also be made as a three-layer system, the bearing lining being made of steel, the actual bearing being made of the alloy of the invention, and the running layer being made of tin or a Sn-rich coating. The multilayer system has a particularly advantageous effect on the suitability and ease of running-in of the sliding bearing and improves the embeddability of foreign and abrasive particles without damage caused by the overlapping of the layer composite system due to pore formation and crack formation in the boundary region of the individual layers even under thermal or thermomechanical conditions on the sliding bearing.
By using the alloys according to the invention, the great potential of copper-nickel-tin materials can be exploited for tin-plated components, wire elements, lead elements and connection elements in electronic and electrical engineering, in particular in terms of strength, elasticity and resistance to stress relaxation. The microstructure of the invention thus reduces the failure mechanism of pore formation and crack formation in the boundary region between the alloy of the invention and tin plating, even at elevated temperatures, which counteracts any increase in the electrical channel resistance of the component or even the separation of the tin plating.
The machining of semifinished products and components made of conventional copper-nickel-tin kneaded alloys, which in each case have a Ni content and a Sn content of up to about 10% by weight, is very difficult to achieve due to insufficient removability of the material. The occurrence of long chips in particular therefore leads to long machine downtimes, since the chips must first be removed by hand from the machining zone of the machine.
In contrast, in the alloy of the present invention, different hard particles act as chip breakers. The short brittle chips and/or tangled chips thus produced contribute to the removability of the material, and therefore the semi-finished products and components made from the as-cast condition and the further worked condition of the alloy of the invention have better machinability.
Tables 1 to 10 illustrate an important processing example of the present invention. The cast slabs of the copper-nickel-tin alloy of the invention and of the reference material are produced by continuous casting. The chemical composition of the castings can be seen from table 1.
Table 1 shows the chemical composition of processing example a and reference material R. Processing example a is characterized by a Ni content of 6.0 wt%, a Sn content of 5.75 wt%, a Si content of 0.3 wt%, a B content of 0.15 wt%, a P content of 0.070 wt% and the balance copper. The reference material R, a conventional copper-nickel-tin-phosphorus alloy, has a Ni content of 5.78% by weight, a Sn content of 5.75% by weight, a P content of 0.032% by weight, and the balance copper.
Table 1: processing of chemical composition (in%) of example A and reference Material R
Alloy (I) Cu Ni Sn Si B P
A Balance of 6.0 5.75 0.3 0.5 0.070
R Balance of 5.78 5.75 - - 0.032
The microstructure of the continuously cast slab of reference material R has pores and shrinkage cavities and Sn-rich segregation especially at grain boundaries.
The continuous cast piece of working example a had a microstructure of uniform solidification, non-porosity and non-segregation due to the effect of the crystallization seeds, compared to reference material R.
The as-cast metal-based material of processing example a consisted of a solid copper solution with about 10 to 15 volume percent, based on the overall microstructure, of an intergrown first phase composition in the form of islands that can pass the empirical formula CuhNikSnmReported, and has a ratio of element contents (h + k)/m in 2 to 6 atomic%. Can be used forDetection of the Compound CuNi having a ratio (h + k)/m of 3.4 and 414Sn23And CuNi9Sn20. Intergrowth in the form of islands in the metal-based material, about 5 to 10 volume percent, based on the overall microstructure, of a second phase component that can pass through the empirical formula CupNirSnsReported, and has a ratio of element contents (p + r)/s in 10 to 15 atomic%. Compound CuNi3Sn8And CuNi4Sn7It was determined to have a ratio (p + r)/s of 11.5 and 13.3. The first phase component and the second phase component of the metal-based material crystallize and encapsulate primarily in the region of the crystallization seeds.
Analysis of first type hard particles in cast State of working example A Compound SiB as represented by Si-containing phase and B-containing phase6For Ni represented by Ni-Si boride6Si2B. For Ni as represented by Ni boride3B. For Ni as represented by Ni phosphide3P and Ni as representative of Ni silicide2Si gives an indication that these compounds are present alone in the microstructure and/or as addition compounds and/or mixed compounds. Furthermore, these hard particles are encapsulated by the first phase component and/or the second phase component of the tin and/or metal based material.
During the casting of working example a, a substructure was formed in the primary cast grain. The grain size of these subgrains in the cast microstructure of inventive process example a was less than 10 μm. Due to the precipitated sub-grain structure and hard particles in the microstructure of inventive processing example A, the as-cast hardness HB at 156 is much higher than the hardness of 94HB of the R continuous cast part (Table 2).
Table 2: hardness HB2.5/62.5 for the as-cast and alloy A and R temper aged at 400 deg.C/3 h/air
Figure GDA0002714863460000221
Figure GDA0002714863460000231
Also shown in table 2 are the hardness values determined on the continuous castings of age-hardened alloys a and R at 400 c over a3 hour duration. At its maximum the reference material R has a hardness of 94 to 145 HB. The hardening is due in particular to the thermally activated formation of Sn-rich phase segregation in the microstructure. The tin-rich phase component precipitates out in a finer form in the hard particle region in the microstructure of working example a. Thus, the increase in hardness from 156HB to 176HB is not significant.
It is an object of the present invention to maintain the good cold formability of conventional copper-nickel-tin alloys despite the introduction of hard particles. To verify the extent to which this goal was achieved, manufacturing procedure 1 was performed according to table 3. The manufacturing procedure consists of one cycle of cold forming and annealing operations, wherein the cold rolling steps are each carried out with the greatest possible degree of cold forming.
Due to the high hardness of the cast state of working example a, it was calcined at a temperature of 740 ℃ for 2 hours and then cooled in an accelerated manner in water. This leads to the assimilation of the properties of the cast state of a and R with respect to strength and hardness.
The achievable cold formability of 60% and 91% for processing example a underscores the fact that the alloy of the invention, despite containing hard particles, can achieve shape change properties even exceeding that of the conventional copper nickel tin alloy R. .
The thermal sensitivity of the reference material R to Sn-rich segregation formation was also found in the annealing (No. 4 in table 3) between the two cold forming steps. Therefore, the annealing temperature of 740 ℃ for the intermediate annealing of the cold-rolled sheet of alloy a must be reduced to 690 ℃ for R.
Table 3: procedure 1 for the production of a strip made of continuously cast sheet from working example A and reference Material R
Figure GDA0002714863460000232
Figure GDA0002714863460000241
After manufacturing procedure 1 was performed, the strip indices of materials a and R were determined after the last cold rolling operation listed in table 4 and at the time of age hardening completion.
It is evident that the strength and hardness of the cold rolled and age hardened strip of working example a at 300 ℃ are higher than the corresponding properties of the reference material strip R.
With the benefit of the high content of hard particles, recrystallization of the microstructure of alloy a occurs at temperatures above about 400 ℃. This recrystallization causes a decrease in strength and hardness, and therefore the effects of precipitation hardening and spinodal line segregation cannot be exhibited.
After age hardening at 450 ℃, the microstructure of further processed working example a includes a second type of hard particles (labeled 3 in fig. 3).
In addition, more phases precipitated in the microstructure of further processed alloy a. These include the continuous precipitates of the (Cu, Ni) -Sn system, identified as 4 in fig.3, and the hard particles of the third category.
A third type of hard particles having a size of less than 3 μm is a feature of the further processed alloy of the present invention. For working example a, a further working according to the invention, it is actually less than 1 μm (marked 5 in fig. 4) after age hardening at 450 ℃.
Table 4: grain size, electrical conductivity and mechanical index of cold rolled and age hardened strip of alloys A and R after manufacturing procedure 1 (Table 3)
■ not yet completely recrystallized
Figure GDA0002714863460000242
Figure GDA0002714863460000251
Further manufacturing procedures were performed in order to reduce the influence of cold formability and recrystallization temperature on the properties of each alloy. The manufacturing procedure 2 pursued the aim of working the continuously cast sheets of materials a and R by cold forming and annealing operations to obtain strips, using the same parameters for each case of the cold forming degree and annealing temperature (table 5).
Due to the high hardness of the as-cast state of working example a, it was calcined again at a temperature of 740 ℃ for a duration of 2 hours before the first cold rolling step and then cooled in an accelerated manner in water.
Table 5: procedure 2 for the production of a strip made of continuously cast sheet from working example A and reference Material R
Figure GDA0002714863460000252
After the last cold rolling step to a final thickness of 3.0mm, the strip of working example a had the highest strength and hardness values (table 6).
Age hardening at 400 ℃ for 3 hours, strength R due to spinodal segregation of the microstructurem(from 498MPa to 717MPa) and Rp0.2The increase in hardness HB (from 439MPa to 649MPa) and the increase in hardness HB (from 166MPa to 230MPa) are most clear for alloy R. However, the age-hardened microstructure of alloy R is very heterogeneous with grain sizes between 5 and 30 μm. Furthermore, the microstructure of the age-hardened state of the reference material R is marked by discrete precipitates of the (Cu, Ni) -Sn system (reference 1 in fig.1 and 2). Ni phosphide (labeled 2 in fig.1 and 2) is also present in the microstructure of the reference material R in the further processed state.
In contrast, the microstructure of the age hardened strip of inventive process example a was very uniform with a grain size of 2 to 8 μm. Moreover, the structure of working example a lacks discrete precipitates even after age hardening for 3 hours at 450 ℃ and then air cooling. In contrast, the second type of hard particles are detectable in the microstructure. These phases are labeled 3 in fig.5 and 6.
In addition, further phases have precipitated in the microstructure of the further processed alloy a. These include the continuous precipitate of the (Cu, Ni) -Sn system labeled 4 in fig.5 and a third class of hard particles. For working example a, a further working according to the invention, the size of the hard particles of the third type after age hardening at 450 c was even less than 1 μm (marked 5 in fig. 6).
Strength R of strip of alloy A after 400 ℃/3 h/air age hardening due to spinodal segregation of the microstructuremAnd Rp0.2Values of 675MPa and 600MPa are presented. Thus, RmAnd Rp0.2Lower than the index of the corresponding age-hardened state of alloy R. Higher proportions of the alloying element nickel may be added to the alloys of the invention if the strength level of R is a particular requirement.
Table 6: grain size, electrical conductivity and mechanical index of cold rolled and age hardened strip of alloys A and R after manufacturing procedure 2 (Table 5)
■ are heterogeneous
Figure GDA0002714863460000261
The next step involved the test of the hot formability of the alloy a and R continuous castings. For this purpose, the cast sheet was hot rolled at a temperature of 720 ℃ (table 7). For the further processing steps of cold forming and intermediate annealing, the parameters of manufacturing procedure 2 were used.
Table 7: procedure 3 for the production of a strip made of continuously cast sheet from working example A and reference Material R
Figure GDA0002714863460000262
Figure GDA0002714863460000271
During hot rolling of cast plates of reference alloy R, deep thermal cracks are formed even after several passes, which leads to failure of the plates by cracking.
In contrast, the cast sheet of processing example a of the present invention was hot-rolled without damage, and could be manufactured to a final thickness of 3.0mm after a plurality of cold rolling processes and calcination processes. The properties of the age-hardened strip (table 8) correspond mainly to those of the strip produced by the production process 2 (label 6) without thermoforming.
Also comparable is the microstructure of a strip made from working example a of the alloy of the invention, which strip was made with and without a hot forming step. Thus, FIGS. 7 and 8 show the uniform structure of the strip produced by working example A, which was produced by a hot forming stage and subsequent age hardening operation at 400 deg.C/3 h/air cooling. In fig.7 and 8, the second type of hard particles, labeled 3, are again apparent.
In addition, fig.7 shows a continuous precipitate of the (Cu, Ni) -Sn system, labeled 4, and a third type of hard particles. In the microstructure of the further processed variant of working example a, the hard particles of the third type actually exhibit a size (marked 5 in fig. 8) of less than 1 μm.
Analysis of the second and third types of hard particles in this further processing state of processing example a for the compound SiB represented as the Si-containing phase and the B-containing phase6For Ni represented by Ni-Si boride6Si2B. For Ni as represented by Ni boride3B. For Ni as represented by Ni phosphide3P and Ni as representative of Ni silicide2Si again gives an indication that these compounds are present alone and/or as addition compounds and/or mixed compounds in the microstructure. Furthermore, these hard particles are coated with precipitates of the (Cu, Ni) -Sn system.
Table 8: grain size, electrical conductivity and mechanical index of cold rolled and age hardened strip of alloy A after manufacturing procedure 3 (Table 7)
Figure GDA0002714863460000272
Figure GDA0002714863460000281
In the construction of facilities, devices, engines and machinery, many applications require components having relatively high dimensions. This is often the case, for example, in the field of sliding bearings. The production of the respective components requires precursor materials having a suitably large size. Due to the limited producibility of infinitely large castings, it is necessary, if possible, to establish the required material properties also by a small degree of cold forming.
Table 9 lists the process steps used during manufacturing procedure 4. The manufacturing operation is carried out by one cycle of cold forming and annealing operations. Due to the temperature sensitivity determined in conventional continuous casting of relatively high strength and hardness for reference material R and the as-cast condition of working example a, only the cast sheet of alloy a was calcined prior to the first cold rolling operation at 740 ℃.
The first cold rolling operation on the cast sheet of alloy R and the annealed cast sheet of alloy a was carried out with a degree of forming of 16%. An annealing operation was performed at 690 c, followed by a cold rolling operation at 12%. Finally, the age hardening of the strip occurs at temperatures of 350 ℃, 400 ℃ and 450 ℃.
Table 9: production Process 4
Figure GDA0002714863460000282
The low degree of cold forming in the first cold rolling step of 16% together with the subsequent annealing operation at 690 ℃ is not sufficient to eliminate the dendritic and coarse-grained microstructure of the reference material R. Moreover, this thermomechanical treatment enhances the grain boundary coverage of alloy R with Sn-rich segregation.
Cracks form during the second cold rolling step from the surface of the strip deep into the interior of the strip across the grain boundaries of R covered by Sn-rich segregation and across the dendritic structure.
The crack-free and homogeneous microstructure of the strip of process example a was characterized by an arrangement of hard particles of the second type and hard particles of the third type. The size of the hard particles of the third type is less than 1 μm even after this manufacturing process 4, as is the case after the foregoing manufacturing process.
The properties of the resulting strip after the last cold rolling operation and after the age hardening operation are shown in table 10. Due to the high density of cracks, it is not possible to take undamaged tensile samples from the strip of material R. Therefore, metallographic analysis and hardness measurements can be performed only on these strips.
Working example a had a high age hardenability as demonstrated by the interaction of the precipitation hardening of the microstructure and the mechanism of spinodal line segregation. Thus, the index R is due to age hardening at 400 ℃ from 517MPa to 639MPa and 481MPa to 568MPamAnd Rp0.2And (4) increasing.
Table 10: after manufacturing procedure 4 (Table 9), the grain size, electrical conductivity and mechanical index of the cold rolled and age hardened strip of alloys A and R
■ dendritic with Sn-rich segregation
Figure GDA0002714863460000291
Thus, it can be shown that the degree of precipitation hardening and spinodal line segregation of the microstructure of the present invention can be tuned to achieve desired material properties through variations in the chemical composition, the degree of forming in cold forming operations, and through variations in the age hardening conditions. In this way, the strength, hardness, ductility and electrical conductivity of the alloy according to the invention can be made in particular in accordance with the envisaged field of use.
List of reference numerals
Discontinuous precipitates of the 1 (Cu, Ni) -Sn system
2 Ni phosphide
3 hard particles of the second type
Continuous precipitates of the 4 (Cu, Ni) -Sn system and hard particles of the third type
5 hard particles of the third type

Claims (16)

1. A high strength copper-nickel-tin alloy in a as-cast condition having excellent castability, hot formability and cold formability, high abrasion resistance to abrasive wear, high galling wear resistance, high fretting wear resistance, and improved corrosion resistance and stress relaxation resistance consisting of, in weight%:
2.0% to 10.0% of Ni,
2.0 to 10.0% of Sn,
0.1 to 0.6% of Si,
0.002% to 0.45% of B,
0.001% to 0.09% of P,
alternatively, up to a maximum of 2.0% Co,
alternatively, up to a maximum of 2.0% Zn,
alternatively, up to a maximum of 0.25% Pb,
and the balance: copper and unavoidable impurities in the form of copper,
the method is characterized in that:
-the Si/B ratio of the element content in% by weight of the elements silicon and boron has a minimum value of 0.4 and a maximum value of 8;
the following microstructural elements are present in the alloy after casting:
a) a Si-containing and P-containing metal-based component having, based on the overall microstructure:
a1) by empirical formula CuhNikSnmReported, and having a first phase composition of up to 35% by volume of the (h + k)/m ratio of the element content in the range of 2 to 6 atomic%,
a2) by empirical formula CupNirSnsReported, and having a second phase composition of up to 15% by volume of the (p + r)/s ratio of the element content in the range of 10-15 atomic%, and
a3) the balance of the solid copper solution;
b) based on the phases present in the overall microstructure,
b1) as Si-containing and B-containing phases in the form of silicon boride and borosilicate and/or borophosphosilicate in an amount of 0.01 to 10% by volume,
b2) based on 1% by volume to 15% by volume, as Ni having the empirical formulaxSi2B, 4 to 6 of Ni-Si boride,
b3) based on 1 to 15% by volume, as Ni boride,
b4) as Ni phosphide in an amount of 1 to 5% by volume,
b5) as Ni-silicides in the microstructure, present alone and/or as addition compounds and/or as mixed compounds, and surrounded by tin and/or a first phase component and/or a second phase component, in a volume of 1% to 5%;
during casting, the Si-containing phase and the B-containing phase, Ni-Si boride and Ni boride, Ni phosphide and Ni silicide in the form of silicon boride, present alone and/or as addition compounds and/or mixed compounds, constitute seeds for uniform crystallization during solidification/cooling of the melt, so that the first phase component and/or the second phase component are distributed uniformly in an island-like and/or network-like microstructure;
the Si-containing phase and the B-containing phase in the form of borosilicate and/or borophosphosilicate together with the phosphosilicate act as a wear-and corrosion-protective coating on the semifinished materials and components of the alloy.
2. A high strength copper-nickel-tin alloy after further treatment by an annealing operation or by a hot forming operation and/or a cold forming operation and an annealing operation, having excellent castability, hot formability and cold formability, high abrasive wear resistance, high adhesive wear resistance and high fretting resistance and improved corrosion resistance and stress relaxation resistance, consisting of, in% by weight:
2.0% to 10.0% of Ni,
2.0 to 10.0% of Sn,
0.1 to 0.6% of Si,
0.002% to 0.45% of B,
0.001% to 0.09% of P,
alternatively, up to a maximum of 2.0% Co,
alternatively, up to a maximum of 2.0% Zn,
alternatively, up to a maximum of 0.25% Pb,
and the balance: copper and unavoidable impurities in the form of copper,
the method is characterized in that:
-the Si/B ratio of the element content in% by weight of the elements silicon and boron has a minimum value of 0.4 and a maximum value of 8;
-after further treatment of the alloy by at least one annealing operation or by at least one hot and/or cold forming operation and at least one annealing operation, the following microstructural composition is present:
A) a metal-based component having, based on an overall microstructure:
A1) by empirical formula CuhNikSnmReported, and having a first phase composition of up to 15% by volume of the (h + k)/m ratio of the element content in atomic% of 2 to 6,
A2) by empirical formula CupNirSnsReported, and having a second phase composition of up to 5% by volume of the (p + r)/s ratio of the element content in the range of 10-15 atomic%,
A3) the balance of the solid copper solution;
B) based on the phases present in the overall microstructure,
B1) the Si-containing phase and the B-containing phase in the microstructure as silicon boride and borosilicate and/or borophosphosilicate in 2 to 30% by volume have the empirical formula NixSi2Ni-Si borides of the 4 to 6 x type, as Ni borides, Ni phosphides and Ni silicides, individually and/or as addition compounds and/or mixed compounds, and encapsulated by a precipitate of the (Cu, Ni) -Sn system,
B2) as a continuous precipitate of the (Cu, Ni) -Sn system in the microstructure in amounts of up to 80% by volume,
B3) as Ni phosphide and Ni silicide in the microstructure in 2 to 30% by volume, these compounds being present individually and/or as addition compounds and/or mixed compounds, being encapsulated by precipitates of the (Cu, Ni) -Sn system and having a size of less than 3 μm;
the Si-containing phase and the B-containing phase, Ni-Si boride and Ni boride, Ni phosphide and Ni silicide in the form of silicon borides present alone and/or as the addition compound and/or the mixing compound constitute seeds for static and dynamic recrystallization of the microstructure during further processing of the alloy, which enables the establishment of a homogeneous and fine-grained microstructure;
the Si-containing phase and the B-containing phase in the form of borosilicate and/or borophosphosilicate together with the phosphosilicate act as a wear-and corrosion-protective coating on the semifinished materials and components of the alloy.
3. The copper-nickel-tin alloy of claim 1 or 2, wherein the elemental nickel and tin are each present in a range of 3.0 wt% to 9.0 wt%.
4. The copper-nickel-tin alloy of claim 1 or 2, wherein the elemental boron is present in an amount of 0.01 to 0.4% by weight.
5. The copper-nickel-tin alloy of claim 1 or 2, wherein the elemental phosphorus is present in an amount of 0.01 wt% to 0.09 wt%.
6. The copper-nickel-tin alloy of claim 1 or 2, wherein the alloy is free of lead except for any unavoidable impurities.
7. Method for producing a component in the form of a final or near-final product from a copper-nickel-tin alloy according to any one of claims 1 to 6 with the aid of a sand casting process, a shell mould casting process, a precision casting process, a full mould casting process, a pressure die casting process or a lost foam process.
8. Method for producing strips, sheets, plates, bolts, round wire, profile wire, round bar, profile bar, hollow bar, tubes and profiles from the copper-nickel-tin alloy according to any one of claims 1 to 6 with the aid of a permanent die casting process or a continuous or semi-continuous casting process.
9. The method of claim 8, wherein the further processing of the as-cast condition comprises at least one hot forming operation at a temperature in the range of 600 ℃ to 880 ℃.
10. The method according to any one of claims 7 to 9, characterized in that at least one annealing treatment is carried out at a temperature ranging from 170 ℃ to 880 ℃ for a time ranging from 10 minutes to 6 hours.
11. The method of claim 8, wherein the further processing of the as-cast condition, the as-hot-formed condition, the as-annealed-as-cast condition, or the as-annealed-as-hot-formed condition comprises performing at least one cold forming operation.
12. The method according to claim 11, characterized in that the annealing treatment is carried out at least once at a temperature ranging from 170 ℃ to 880 ℃ for a time duration ranging from 10 minutes to 6 hours.
13. Method according to claim 11 or 12, characterized in that the stress relief annealing/age annealing operation is carried out at a temperature in the range of 170 ℃ to 550 ℃ for a time of 0.5 to 8 hours.
14. Use of the copper-nickel-tin alloy according to any one of claims 1 to 6 for movable and sliding clamping strips, for friction rings and discs, for sliding and guiding elements in internal combustion engines, valves, turbochargers, gears, exhaust gas aftertreatment systems, lever systems, brake systems and joint systems, machines and equipment in hydraulic units or in general mechanical engineering.
15. Use of the copper-nickel-tin alloy according to any one of claims 1 to 6 for components, wire elements, guiding elements and connecting elements in electronic/electrical engineering.
16. Use of the copper-nickel-tin alloy of any one of claims 1 to 6 for metal articles in breeding of marine habitats, for percussion instruments, for propellers, wings, marine propellers and hubs for shipbuilding, for housings for water, oil and fuel pumps, for guide, runner and paddle wheels for pumps and water turbines, for gears, worm gears, helical gears, gland and spindle nuts, and for pipes, seals and connecting bolts in the marine and chemical industries.
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